Precipitation hardening nickel-base alloy, part made of said alloy, and manufacturing method thereof

ABSTRACT

A precipitation hardened nickel-base alloy, characterized in that its composition is, in weight percentages:
         18%≦Cr≦22%, preferably 18%≦Cr≦20%;   18%≦Co≦22%, preferably 19%≦Co≦21%;   4%≦Mo+W≦8%, preferably 5.5%≦Mo+W≦7.5%;   trace amounts≦Zr≦0.06%;   trace amounts≦B≦0.03%. preferably trace amounts≦B≦0.01%;   trace amounts≦C≦0.1%, preferably trace amounts≦C≦0.06%;   trace amounts≦Fe≦1%;   trace amounts≦Nb≦0.01%;   trace amounts≦Ta≦0.01%;   trace amounts≦S≦0.008%;   trace amounts≦P≦0.015%;   trace amounts≦Mn≦0.3%;   trace amounts≦Si≦0.15%;   trace amounts≦O≦0.0025%;   trace amounts≦N≦0.0030%;       

     the remainder being nickel and impurities resulting from the elaboration, the Al and Ti contents further satisfying the conditions: 
       Ti/Al≦3;  (1)
 
       Al+1.2 Ti≧2%;  (2)
 
       (0.2 Al−1.25) 2 −0.5 Ti≧0%;  (3)
 
       Ti+1.5 Al≦4.5%.  (4)
 
     Part made in this alloy and its manufacturing method.

The invention relates to nickel-base alloys (superalloys), and more specifically those intended for manufacturing the parts which have to be used at high temperatures. Typically, this is the case of elements of land, aeronautical and other turbine engine elements.

For these type of uses, an alloy NiCo20Cr20MoTi (AFNOR standard) so called “C263” alloy is known, the composition of which is typically Ni, Cr (19-21%), Co (19-21%), Mo (5.6-6.1%), Ti (1.9-2.4%), Al (<0.6%). The percentages are weight percentages as this will be the case for all the compositions indicated subsequently.

This is a precipitation hardening alloy, the latter being ensured by the presence of the γ′ phase (Ni₃Ti, Al), and which has good forgeability and weldability properties. As regards this last point, this is due to that, unlike what is often encountered for alloys hardened by the γ′ phase, it is not subject to the cracking phenomenon due to fragilization by strain age cracking in the welding areas. It also has good hot tensile ductility and satisfactory hot strength. Generally, its weldability/forgeability compromise is advantageous.

However, it has the drawback of having microstructural instability between 700 and 900° C., a range of temperatures in which the phase η may be formed to the detriment of the γ′ phase (see the reference: Zhao, Metallurgical and Materials Transactions A, 2001, Vol. 32A, pp 1271-1282). The ductility and the resilience are thereby degraded. It is therefore not optimally adapted to uses where the parts are brought to such temperatures.

Other alloys are known for such uses and do not have this structural instability, but they have other drawbacks.

The alloy known under the name of INCO 617 (Ni, Cr (20-24%), Co (10-15%), Mo (8-10%), Al (0.8-1.5%), Ti (0-0.6%)) has a good weldability/forgeability compromise, but its hot mechanical properties (notably at about 750° C. which is a frequent temperature of use for the parts to which the invention preferentially applies) are insufficient.

The alloy known under the name of RENE 41 (Ni, Cr (18-20%), Co (10-12%), Mo (9-10.5%), Al (1.4-1.6%), Ti (3-3.3%)), conversely has good hot mechanical properties but its weldability/forgeability compromise is not optimum. The same applies to the alloy known under the name of WASPALOY (Ni, Cr (18-21%), Co (12-15%), Mo (3.5-5%), Al (1.2-1.6%), Ti (2.75-3.25%). These insufficient weldability/forgeability compromises are probably due to a too great proportion of γ′ phase.

Therefore there exists a need for industrialists to have available Ni-base alloys for uses at high temperature (typically 700-900° C.) having both good microstructural stability at the temperatures of use, good mechanical properties at these same temperatures, and simultaneously good forgeability and good weldability allowing the manufacturing of said parts in the desired configurations and their integration into the devices for which they are intended.

For this purpose, the object of the invention is a nickel-base alloy with precipitation hardening, characterized in that its composition is, in weight percentages:

-   -   18%≦Cr≦22%, preferably 18%≦Cr≦20%;     -   18%≦Co≦22%, preferably 19%≦Co≦21%;     -   4%≦Mo+W≦8%, preferably 5.5%≦Mo+W≦7.5%;     -   trace amounts≦Zr≦0.06%;     -   trace amounts≦B≦0.03%, preferably trace amounts≦B≦0.01%;     -   trace amounts≦C≦0.1%, preferably trace amounts≦C≦0.06%;     -   trace amounts≦Fe≦1%;     -   trace amounts≦Nb≦0.01%;     -   trace amounts≦Ta≦0.01%;     -   trace amounts≦S≦0.008%;     -   trace amounts≦P≦0.015%;     -   trace amounts≦Mn≦0.3%;     -   trace amounts≦Si≦0.15%;     -   trace amounts≦O≦0.0025%;     -   trace amounts≦N≦0.0030%;

the remainder being nickel and impurities resulting from the elaboration, the Al and Ti contents further satisfying the conditions:

Ti/Al≦3;  (1)

Al+1.2 Ti≧2%;  (2)

(0.2 Al−1.25)²−0.5 Ti≧0%;  (3)

Ti+1.5 Al≦4.5%.  (4)

The γ′ phase fraction is preferably comprised between 5 and 20%.

The solvus temperature of its γ′ phase is preferably less than or equal to 980° C.

The object of the invention is also a method for manufacturing a part in a nickel-base alloy, characterized in that an ingot is prepared having the composition defined earlier, it is homogenized at a temperature of at least 1,150° C. for 24 to 72 h, it is hot-worked by forging or rolling in a range of supersolvus temperatures, it is solution treated at a temperature from 1,100 to 1,200° C. for 1 to 4 h, it is cooled at at least 1° C./min, for example in water, it is aged at a temperature from 750 to 850° C. for 7 to 10 h, and it is cooled, for example in calm air or in an enclosure.

The object of the invention is also a part in a nickel-base alloy, characterized in that it was prepared according to the previous method.

For example this is an element of a land or aeronautical turbine engine.

As this will have been understood, the invention is based on an optimization of the known C263 grade, which essentially consists in an equilibrium cleverly selected between the Al and Ti contents. This equilibrium will drive:

-   -   The stability of the γ′ phase at a high temperature (700-900°         C., in particular 750° C.), in order to avoid its transformation         into an acicular phase η (of composition Ni₃Ti, therefore         without any Al);     -   The γ′ phase fraction formed at 700-900° C., in particular at         750° C.;     -   The solvus temperature of the γ′ phase.

On the remainder of the composition of the alloy, the changes with respect to the known C263 are small, and it was checked that the optimizations of the Al and Ti contents according to the invention do not lead to a modification of the advantageous properties of the alloy which are not directly related to the γ′ phase.

The invention will be better understood by means of the description which follows, given with reference to the following appended figures:

FIGS. 1 to 8 which show micrographs of reference samples (FIGS. 1 and 5 to 8) and according to the invention (FIGS. 2 to 4);

FIG. 9 which shows the results of tests of measurement of the tensile strength Rm of these samples versus temperature;

FIG. 10 which shows the results of measurement tests of the conventional elastic limit Rp_(0.2) of these samples versus temperature;

FIG. 11 which shows the results of measurement tests of the elongation at break A % of these samples versus temperature;

FIG. 12 which shows the results of tests of measurement of striction Z % of these samples versus temperature;

FIG. 13 which shows the results of creep tests at failure at 750° C. of these samples, wherein the breakage stress is given according to the Larson-Miller parameter;

FIG. 14 which shows the results of resilience tests of two samples (a reference sample and a sample according to the invention), conducted after the final heat treatment of the sample and after over-ageing at 750° C. for 3,000 h representative of what the metal may undergo during a use for which it is preferentially intended;

FIGS. 15 to 18 which show a sample according to the invention and reference samples during forging.

A first condition for the optimization of the equilibrium between Al and Ti is that the formation of a η phase is avoided at the temperatures of use of the alloy during its preferential uses, i.e. at temperatures of 700-900° C., typically of the order of 750° C. The formation of the η phase is directly related to the Ti and Al contents present in the alloy and to their ratio. The ranges of contents of these elements therefore have to be determined, which allow them to be avoided at 700-900° C., considering the reminder of the composition of the alloy. Thermodynamic calculations, carried out by means of the THERMOCALC software currently used by metallurgists and which was also used in a first approach for the remainder of the optimization, have indicated that for C263, if the Ti/Al ratio is less than or equal to 3, the formation of an η phase was avoided, and this regardless of the Al level in the alloy.

The condition must therefore be observed:

Ti/Al≦3  (1)

Another condition is that for guaranteeing the tensile strength and creep strength properties at 700-900° C., the atomic percentage of the phase γ′ present at these temperatures in the alloy should be of at least 5%. Below this value, one does not have sufficient precipitation hardening. It is estimated that this condition is fulfilled when the weight percentages of Al and Ti observe the relationship:

Al+1.2 Ti≧2%.  (2)

As regards the forgeability (or hot deformability generally speaking, for example by rolling) and weldability properties, it is possible to state the following.

Under standard forging conditions at a high temperature, the forging is carried out in a temperature range where there is no precipitation of the γ′ phase which would make the metal too hard and subject to the occurrence of defects, such as cracks, during deformations. It is therefore carried out at a temperature greater than the solvus temperature of this phase. This temperature therefore has the advantage of not being too high, so forging is possible under industrial conditions. More specifically, the solvus temperature of the γ′ phase should be as low as possible in order to avoid precipitation of this phase during the inevitable cooling of the product during forging.

The fraction of the γ′ phase which may precipitate at high temperature should also be taken into account. Indeed, the greater the fraction of the precipitated hardening phase at a high temperature, the more the alloy may harden during variations of temperatures which may occur during forging, which may complicate the execution of the operation. This undesired precipitation of γ′ phase at this specific moment of the preparation of the product also has influence on the weldability, because of the possibility of cracking due to fragilization under hot stress. Indeed, the larger the fraction of precipitated γ′ phase in the welded area, the higher are the stresses generated by the precipitation of the γ′ phase in this same area during cooling and they promote therein cracking after the welding.

In order that the good conditions required for hot formability and weldability be simultaneously satisfied, it is therefore necessary to retain a solvus temperature of the γ′ phase of at most 980° C., and to limit the γ′ phase fraction present at 700-900° C. to 20% (in atomic %), in particular at 750° C.

These conditions are observed if the weight contents of Ti and Al observe both conditions:

(0.2 Al−1.25)²−0.5 Ti≧0%;  (3)

Ti+1.5 Al≦4.5%  (4)

As regards the other elements which have to or may be present, i.e. like mandatory or optional alloy elements, either as impurities to be limited, the following may be stated. The preferred ranges are those where one is most sure of obtaining the mentioned advantages of each element without having the drawbacks thereof.

The Cr content is comprised between 18 and 22%, preferably 18 to 20%. Cr is important for guaranteeing resistance to corrosion and to oxidation, and for establishing the resistance of the alloy to the effects of the high temperature environment. A too high content promotes the obtaining of undesirable fragile phases, such as the σ phase, and the limit of 22% by weight is set accordingly.

The Co content is comprised between 18 and 22%, preferably 19 to 21%. A high Co content is necessary in order to improve the forgeability of the grade by decreasing the solvus temperature of the phase γ′, nevertheless it should be limited, mainly for reasons of cost.

The sum of the Mo and W contents should be comprised between 4 and 8%, preferably 5.5 to 7.5%. Both of these elements may be substituted for each other. The lower limit of 4% guarantees precipitation hardening and good creep strength, and the upper limit of 8% avoids the formation of harmful phases.

The Zr content is comprised between trace amounts (in other words an absence of voluntary addition, the optional residual Zr content only resulting from the melting of the raw materials and of the elaboration, with the associated impurities) and 0.06%.

The B content is comprised between trace amounts and 0.03%, preferably 0.003 to 0.01%.

The C content is comprised between trace amounts and 0.1%, preferably 0.04 to 0.06%.

These last three elements form segregations at the grain boundaries which contribute to the hot strength and ductility by trapping the optionally present noxious elements, like S. They promote creep resistance under conditions of low stresses and of high temperatures. However, if they are present in excess, they reduce the melting temperature of the segregated areas and strongly alter forgeability. Their optional presence should therefore be well controlled.

It should be understood that the preferential contents of the elements which have just been mentioned are independent of each other. In other words, an alloy which would have a preferential content on one or several of them only, but not on the others, should nevertheless be considered as an advantageous embodiment of the invention.

As regards the elements, the contents of which have the advantage of being minimized as much as possible, the following may be stated.

The Fe content is limited to a maximum of 1%. Beyond, it risks forming phases which are detrimental to the properties of the alloy.

The Nb and Ta contents are both limited to a maximum of 0.01%. These elements are expensive and have a great tendency to segregate without these segregations having any advantages which may compensate their drawbacks (unlike what may occur for Zr, B and C).

The S, P, Mn and Si contents should also be limited so as not to reduce hot ductility. An excess of Si would also cause a precipitation of Laves phases during solidification, and it will be difficult to put them back into solution during subsequent heat treatments. Resilience would also be found to be degraded.

The maximum contents accepted for these elements are therefore 0.008% for S, 0.015% for P, 0.3% for Mn, and 0.15% for Si.

In order to guarantee good mechanical properties of the alloy, the 0 content should be limited to 25 ppm at most and the N content to 30 ppm at most. For this purpose, an elaboration in vacuo and also involving a method such as re-melting under an electroconductive slag (ESR) or arc re-melting in vacuo (VAR) is particularly recommended. But from these points of view, the alloys of the invention are not particularly distinguished from the usual C263 grades for which they are intended to be substituted.

As regards the method for preparation of the parts, typically an ingot is prepared having the previous composition by dual or triple melting, therefore by involving at least one of the ESR and VAR methods, it is homogenized at a temperature of at least 1,150° C. for 24 to 72 h, it is hot worked by forging or rolling in a range of supersolvus temperatures, it is solution heat treated at a temperature from 1,100 to 1,200° C. for 1 to 4 h, it is rapidly cooled at at least 1° C./min, for example in water, it is aged at a temperature from 750 to 850° C. for 7 to 10 h, and it is cooled for example in calm air, or in an enclosure. Depending on the targeted applications, alternatives may be provided to this method, by not executing some of these steps or by adding other steps. They may notably be followed by machining or by any other operation for setting the definitive dimensions of the part.

An elaboration of the part resorting to a powder metallurgy method and resulting in a product having the required composition properties would also be conceivable.

Tests were conducted on samples, the compositions of which are mentioned in table 1.

TABLE 1 Compositions of the tested samples O N Samp.. Ni % Cr % Co % Mo % W % B % C % Zr % Al % Ti % ppm ppm (1) (2) (3) (4) A 51.60 19.71 20.15 5.98 Trace 0.005 0.051 0.02 0.77 1.50 3.5 17 1.06 2.57 0.45 2.66 amounts B 47.50 20.86 20.49 5.96 1.43 0.010 0.050 0.02 1.95 1.13 3.1 18 0.58 3.31 0.17 4.06 C 51.00 19.79 20.12 6.13 Trace 0.010 0.050 0.01 2.64 0.22 3.4 15 0.08 2.90 0.41 4.18 amounts D 51.50 19.74 20.00 6.20 Trace Trace 0.052 0.01 0.42 2.24 3.1 22 5.33 3.11 0.24 2.87 amounts amounts E 50.40 19.60 20.00 5.97 Trace 0.002 0.049 0.003 3.00 0.252 3 16 0.08 3.30 0.30 4.75 amounts F 48.20 19.52 20.60 4.22 3.48 0.010 0.050 0.02 3.62 0.15 4 17 0.04 3.80 0.20 5.58 G 49.70 19.97 18.50 7.50 Trace 0.010 0.060 0.02 2.20 1.95 3.3 14 0.89 4.54 −0.32 5.25 amounts H 52.10 20.00 18.20 8.00 Trace Trace 0.060 0.01 1.10 0.48 3.2 16 0.44 1.68 0.82 2.13 amounts amounts

The samples A, B and C correspond to the invention, the other samples are reference alloys which do not observe at least one of the conditions (1) to (4) defined earlier because of their Al and Ti contents. The sample B corresponds to the version of the invention considered as optimum, wherein the contents of all the elements are in the preferred ranges. The reference sample D corresponds to a standard alloy of the C263 type which does not observe the relationship (1). The sample E and the sample F do not observe the relationship (3). The sample G does not observe the relationships (3) and (4). The sample H does not observe the relationship (2). This actually shows that observance of all the relationships (1) to (4) is required in order to obtain the desired results.

The tested samples were elaborated by dual melting VIM-VAR (i.e., as this is standard, by melting of the raw materials in a vacuum induction oven, followed by casting and solidification of an electrode, the latter being refined by re-melting in vacuo in an arc oven), in order to obtain 200 kg ingots. This method is currently used for manufacturing ingots intended to form forged or rolled parts with a high inclusion purity and at low contents of residual elements, notably gaseous elements. However it is not mandatorily used for elaborating the alloys of the invention, if the latter are intended for the manufacturing of parts not having very high requirements on these points. In these cases, less complex conventional elaboration methods may be used, in so far that they give the possibility of attaining the indispensable low levels of certain residual elements, notably by suitable selection of the raw materials.

These ingots were homogenized at a temperature above 1,150° C. for 48 h, and then forged into bars with a diameter of 80 mm between 1,200 and 1,050° C.

The examples then underwent the following heat treatments:

-   -   Solution treatment at 1,140° C.+/−10° C. for 2 h, followed by         water quenching;     -   Ageing at 800° C.+/−10° C. for 8 h followed by cooling in air.

This heat treatment is typical of the C263 alloy for its usual applications such as turbine engine elements.

The THERMOCALC software does not provide any occurrence of the η phase for these samples under their treatment conditions, except for the sample D.

In fact, micrographs were made on portions of said samples having been subject to over-ageing at 750° C. for 3,000 h in order to simulate a use of the corresponding alloys at a high temperature. Micrographs made with a field effect electron microscope are shown in FIGS. 1 (sample D), 2 (sample A), 3 (sample B), 4 (sample C), 5 (sample E), 6 (sample F), 7 (sample G) and 8 (sample H).

It is confirmed that only the sample D, representative of a standard C263 alloy includes a significant amount of acicular η phase (as needles). The other samples, notably those of the invention A, B and C, do not have this phase for which the invention notably aimed at avoiding the occurrence during a use at 700-900° C., typically about 750° C.

FIG. 9 shows the results of tensile mechanical tests on these same samples for the measurement of Rm, carried out between room temperature and 800° C. FIG. 10 shows the results of measurement of Rp_(0.2), FIG. 11 shows the results of the measurement of the elongation at break A %, and FIG. 12 shows the results of striction tests Z %, carried out under the same conditions.

It is found that the alloys B and C according to the invention have tensile test results (Rm and Rp_(0.2)) similar to those of the reference alloy D. The tensile test results of the alloy A according to the invention are slightly degraded as compared with those of alloy D, but remain satisfactory. And the hot ductility of the alloy A is the best of all, which may be an advantage for certain uses. The invention therefore actually gives the possibility of satisfactory optimization or retention of all these hot mechanical properties as compared with the reference C263 alloy.

The alloys E, F and G have very good tensile test results, notably under hot conditions, but their loss of hot ductility is very large, which may be ascribed to poor equilibration of the Al and Ti contents.

The alloy H is unsatisfactory at every point of view at high temperatures.

FIG. 13 shows the results of breakage creep tests at 750° C.: the breakage stress in MPa is given versus the Larson-Miller parameter (PLM) in the standard procedure.

The alloys A, B, C according to the invention, and the reference alloys F and G have higher lifetimes at breakage than the one of the reference alloy D. This shows, also from this point of view, that the invention provides improvement in the performances of the alloy D which is the nearest to them. The alloy E has a short lifetime because of its insufficient hot ductility, and the tests were not able to be prolonged beyond a PLM of 23.4. The alloy H is, there again, very clearly unsatisfactory.

FIG. 14 shows the result of resilience tests conducted on several specimens of the alloy A according to the invention and reference alloy D, after solution heat treatment and then ageing as described above on the one hand, after 3,000 hours of overageing at 750° C. subsequently to the preceding heat treatment on the other hand, there again for simulating the time-dependent change in the alloy during use. The results are clear: the resilience Kv is practically unaffected by overageing of the sample A, while it substantially drops for the sample D. This confirms that the formed η phase during use at a high temperature of the standard C263 alloy has a strong embrittling affect and that the invention gives the possibility of finding a remedy to this problem.

Forging tests were also conducted under identical conditions (homogenization at more than 1,150° C. for 48 h and then forging at 1,200° C.-1;050° C. down to a diameter of 80 mm), and FIGS. 15 to 18 show the obtained results.

The alloys A, B and C according to the invention, as well as the reference alloy H were forged without any problems like the alloy D would have been: no crack appeared during the forging. FIG. 15 shows the alloy A during forging at about 1,100° C. and no crack is actually visible. FIG. 16 shows the alloy E during forging at the same temperature, and slight cracks are visible. FIG. 17 shows the alloy F during forging at the same temperature, and the cracks are much deeper than in the previous cases. FIG. 18 shows the alloy G during forging at the same temperature, and there again deep cracks are visible. The good forgeability of the alloys according to the invention is therefore confirmed, and is ascribed to a lower γ′ phase proportion than for the reference samples E, F and G.

A preferred application of the invention is the manufacturing of land and aeronautical turbine engines, but of course it is not exclusive. 

1. A precipitation hardened nickel-base alloy, wherein its composition consists of, in weight percentages: 18%≦Cr≦22%; 18%≦Co≦22%; 4%≦Mo+W≦8%; trace amounts≦Zr≦0.06%; trace amounts≦B≦0.03%; trace amounts≦C≦0.1%, preferably trace amounts≦C≦0.06%; trace amounts≦Fe≦1%; trace amounts≦Nb≦0.01%; trace amounts≦Ta≦0.01%; trace amounts≦S≦0.008%; trace amounts≦P≦0.015%; trace amounts≦Mn≦0.3%; trace amounts≦Si≦0.15%; trace amounts≦O≦0.0025%; trace amounts≦N≦0.0030%; the remainder being nickel and impurities resulting from the elaboration, the Al and Ti contents further satisfying the conditions: Ti/Al≦3;  (1) Al+1.2 Ti≧2%;  (2) (0.2 Al−1.25)²−0.5 Ti≧0%; and  (3) Ti+1.5 Al≦4.5%.  (4)
 2. The alloy according to claim 1, characterized in that its γ′ phase fraction is comprised between 5 and 20%.
 3. The alloy according to claim 1, characterized in that the solvus temperature of its γ′ phase is less than or equal to 980° C.
 4. A method for manufacturing a part in a nickel-base alloy, wherein an ingot is prepared, having the composition according to claim 1, it is homogenized at a temperature of at least 1,150° C. for 24 to 72 h, it is hot-worked by forging or rolling in a supersolvus temperature range, it is put into solution at a temperature from 1,100 to 1,200° C. for 1 to 4 h, it is cooled at a rate of at least 1° C./min, for example in water, it is aged at a temperature from 750 to 850° C. for 7 to 10 h, and it is cooled for example in calm air or in an enclosure.
 5. A nickel-base alloy part, which was prepared according to the method of claim
 4. 6. The part according to claim 5, wherein it is an element of a land or aeronautical turbine engine.
 7. The alloy according to claim 1, wherein 18%≦Cr≦20%.
 8. The alloy according to claim 1, wherein 19%≦Co≦21%.
 9. The alloy according to claim 1, wherein 5.5%≦Mo+W≦7.5%.
 10. The alloy according to claim 1, wherein trace amounts≦B≦0.01%.
 11. The alloy according to claim 1, wherein trace amounts≦C≦0.06%.
 12. The part according to claim 5, wherein it is an element of an aeronautical turbine engine. 